steel more homogeneous. Nevertheless, cast slabs and blooms must be soaked to eliminate as-cast segregations. This process, to the extent it is done, occurs as they are reheated to the ap- propriate temperature for hot working. Soaking dissolves the few percent of residual delta fer- rite that are present on slab solidification. It is important to soak at the highest temperature at which delta ferrite is not a stable phase so that it will dissolve, about 1250 °C (2280 °F) for most austenitic stainless steels.
Soaking at higher temperatures causes ferrite levels to increase, negating the homogenization and causing very poor hot workability. Longer times at temperature than the minimum required for thermal uniformity also cause problems as any sulfur and oxygen impurities are rejected from austenite and can diffuse to grain bound- aries, where they form weak, plastic films that also degrade hot workability. Grain growth, by reducing grain boundary area, exacerbates this effect. Thus, soaking times are best minimized and closely controlled. Alloys are therefore de- signed to have only a slight amount of delta fer- rite to be redissolved during soaking. Ferrite is useful because it has a high solubility for oxygen and sulfur. Having none would result in impurity rejection of these elements to grain boundaries during solidification, which results in the worst- possible hot working characteristics. The oxygen and sulfur trapped in the ferrite during solidifica- tion precipitates in the solid state as inclusions, which also must be equilibrated with the sur-
rounding matrix by sufficient soaking. Welds that are unannealed have such precipitated inclusions in an unequilibrated state, and the result is dimin- ished chromium concentration and poorer pitting resistance.
Annealing
Annealing serves two main functions in stainless steel: It removes the effect of cold work by replacing strained microstructure with new strain-free grains, that is, recrystallization. New grains nucleate and grow. If stored strain energy is insufficient, as happens often with ferritic stainless steels, true recrystallization is difficult to achieve, and the annealing process may only produce recovery without recrystal- lization, leaving the same grains relieved of stress. This leaves the surviving grains with the same crystallographic orientation that deforma- tion produced and may or may not be the de- sired outcome. Second, annealing returns into solution solute that has been precipitated as un- wanted phases, principally carbides, but also intermetallic phases. Annealing also may help to reduce solute segregation remaining from the casting process, making the composition more homogeneous. The homogenization process is accelerated by the reduction in di- mensions from hot and cold working. A reduc- tion in dimension by a factor of two reduces the time to achieve a given degree of homogeniza- tion by a factor of four.
Table 1 Recommended thermal processing temperatures for austenitic alloys
Annealing Annealing ASTM A480(a) Stress Stress
Alloy temperature, °C temperature, °F 2006, °F relieving, °F relieving, oC
Standard alloys . . . 1500–1600 non-L, 815–870 non-L, 1000–1600 L grades 540–870 L grades 201, 202, 201LN 1010–1120 1850–2050 1900 min . . . . 301, 301LN, all versions 1010–1120 1850–2050 1900 min . . . . 304, 304L, 305, all versions 1010–1120 1850–2050 1900 min . . . . 316, 316L, 316N, 317, 317L 1040–1175 1900–2150 1900 min . . . . 308, 309, 309S, 310, 310S 1040–1175 1900–2150 1900 min . . . . Stabilized alloys . . . 1000–1600 540–870 321 955–1065 1750–1950 1900 min . . . . 347, 348 980–1025 1800–1950 1900 min . . . . 20Cb-3 925–955 1700–1750 . . . 925–1010 . . . Moderately alloyed, Creq<26 1120–1175 2050–2150 Various 1500–1600 815–870 S31725, N08028, JS700
Highly alloyed, lower sigma 1120–1175 2050–2250 Various Not recommended . . . alloys, Creq>30, high N
AL6-XN, 4565, 654SMO, 254SMO
Highly alloyed, sigma-prone 1205–1230 2200–2250 Various Not recommended . . . alloys, Creq>30, low N
AL6-X
Annealing to recrystallize is fairly rapid. To a first approximation, it is instantaneous, and the results are merely a function of the maximum temperature attained. This may not be the case for continuous annealing lines, in which transit time can be short enough, less than a minute at temperature, to limit the grain size attained. The driving force for recrystallization is the strain energy stored in the lattice from deformation. The strain energy in a given material is propor- tional to the square of the material’s flow stress. As the material is heated, recovery occurs first. This is the change in physical and mechanical properties associated with dislocation annihila- tion and polygonalization that occurs before the nucleation and growth of new grains.
The nucleation of new grains occurs at high- angle grain boundaries and proceeds by the movement of roughly hemispherical growth fronts into strained areas. The percentage re- crystallized, once a sufficient temperature is reached, grows sigmoidally. It is normal for the time to fully recrystallize to be rather less than the time to attain that temperature. Even at the lower range of annealing temperatures, times are generally less than 1 min. Recrystal- lization will not occur if the lattice contains in- sufficient strain energy. Thirty percent cold work should be used as a rough threshold for the required amount. Annealing after lower amounts of cold work is characterized by scarce nucleation sites and can result in very large and nonuniform grain size. Hot-worked material often has a composite structure that may have already had some recrystallization depending on the final reduction temperature. Annealing may not produce a clear recrystal- lized structure in this case.
The relative rapidity of recrystallization an- nealing is due to the fact that it is rate con- trolled by short-range diffusion. Solution an- nealing requires longer-range diffusion and thus can require much longer times. Some stud- ies have shown that welds, for instance, do not recover completely from their loss of corrosion properties that arise from local alloy depletion unless they have been annealed for times on the order of 1 h (Ref 1). Others have seen homoge- nization in as little as 10 min (Ref 2). Wrought materials can require much shorter times be- cause reductions during hot working have re- duced diffusion distances. It should be noted that precipitates can be redissolved and not ap- parent in the annealed microstructure without full homogeneity being achieved. For example,
carbides can be redissolved and carbon dif- fused away from the carbide, but this does not mean that all composition gradients have been reduced to zero. This means that precipitates may re-form more rapidly in such a material than they would in a completely homogeneous alloy.
The annealing temperature for a given alloy is chosen based on the temperature required to put all alloying elements into solution. Higher car- bon levels, for instance, require higher tempera- tures to dissolve all the carbon. This is one of the principle values of accurate phase diagrams. If it were simply a consideration of recrystal- lization, all alloys could be annealed at similar temperatures at the low end of the recom- mended range. Within the recommended range, the temperature selected should be determined by the desired grain size. End use determines whether a fine or coarse grain size is preferable. Table 1 lists recommended annealing tempera- tures for austenitic stainless steels.
The overall interplay between prior cold work and annealing temperature on mechanical prop- erties of annealed material can be summarized as (Ref 3):
• Grain size of a given alloy is the most im-
portant parameter in characterizing mechan- ical properties.
• Yield and tensile strength are essentially
constant for a given grain size regardless of the amount of prior cold work; however, the elongation depends on the prior reduction. • Yield strength, tensile strength, and hardness
are essentially linear functions of grain size.
• Elongation decreases with finer grain size
and at an increasing rate as grain size be- comes finer as long as cross-section size of the test specimen is not extremely small. This is not true for very coarse-grained ma- terial.
• Maximizing elongation comes from maxi-
mum prior cold work and medium annealing temperatures
• Anisotropy coefficients, or plastic strain ra- tios r are constant up to about 40% reduction after, which r45and rnincrease sharply, while
rt decreases. This leads to earing during
drawing.
• The increase in properties for a one ASTM
grain size increment is:
a. 13 MPa (2 ksi) for tensile strength b. 20 MPa (3 ksi) for yield strength
Atmospheres for annealing are important. Austenitic stainless steels heated in air, of course, form oxide scales. Beneath this oxide, the metal matrix becomes significantly depleted of chromium (Ref 4), often more than 5% lower in chromium and to a depth of as much as 10 µ (395 µin.). So, not only must any oxide be re- moved, so must the chromium-depleted layer. This requires aggressive pickling, which while done commonly, may not be practical for many stainless users. The chromium-depleted zone, however, does pickle rapidly precisely because it does have less chromium. To avoid oxide scale formation, vacuum, hydrogen, or inert gas atmospheres may be used.
If vacuum is used, it should be less than 2 × 10–3torr (0.3 Pa). If an inert gas or hydrogen
is used, the key consideration is moisture con- tent. The dew point must be –40 °C (–40 °F) or lower. More stringent levels may be required if mirror finishes are desired. Cool down must be rapid as oxidation potential increases as tempera- ture decreases. Vacuum or inert gas is preferable to hydrogen for alloys containing stable oxide formers such as aluminum or titanium or for al- loys containing boron.
Austenitic alloys that are subject to sensitiza- tion must be cooled rapidly enough from anneal- ing temperatures to avoid carbide precipitation during cooling. If forced air or water quenching are impractical or if section size prohibits rapid cooling, then using stabilized or low-carbon grades is indicated.
Superaustenitic stainless steels, and even al- loys like 317, present a special problem because these alloys have significant sigma-forming ten- dencies. Sigma forms initially because solidifi- cation segregation causes local enrichment of sigma-promoting elements, such as molybde- num. It can also form from slow cooling of slabs or hot bands. This latter sigma forms at grain boundaries and will cause embrittlement and re- duced corrosion resistance, so it must not only be redissolved, but also the alloy must be ho- mogenized to remove the residual concentration gradients from the sigma. If this is not done, chromium- and molybdenum-depleted regions will still exist, and sigma will re-form much more rapidly during subsequent exposure to high temperatures. For this reason, the higher ends of the annealing ranges are recommended, and annealing times should be generous. Newer alloys have higher nitrogen contents to suppress formation of sigma and other deleterious inter-
metallic phases. Use of high-chromium and- molybdenum alloys without enhanced nitrogen is no longer recommended, and the use of lower- nitrogen alloys should be reexamined and ques- tioned if specified.
Last, stainless surfaces should be scrupu- lously clean before annealing. Even hard water deposits can cause differential oxide growth, which can cause etched spots on the surface, where the postanneal pickling attacks the differ- ent oxide more strongly. Carbonaceous materi- als left on the surface are even more objection- able because they can cause carburization and subsequent loss of corrosion resistance.
Stabilizing anneals are sometimes conducted on stabilized alloys such as 321 and 347. This is useful when carbon levels are sufficiently high that significant dissociation of carbides occurs at annealing temperatures. A second anneal at lower temperature, about 900 °C (1650 °F), then is done to permit the carbon to combine with the stabilizing element rather than leaving it available to form chromium carbides. Current preferred practice for these alloys is to maintain carbon and nitrogen below 0.03% for corrosion- resistant service, which renders this stabilizing unnecessary. Alloys used for high-temperature service benefit from the creep-resisting contri- butions of higher carbon levels.
Stress Relieving
Austenitic stainless steel weldments often con- tain residual stresses, which can cause distortion or lead to stress corrosion cracking in service. They are commonly stress relieved at tempera- tures slightly below the annealing temperature, so that residual stresses may be relieved by creep. One hour at 900 °C (1650 °F) reduces residual stress by about 85%. Lower tempera- tures require exponentially longer times for the same stress relief, with times doubling for each 100 °C (180 °F) decrement as decreasing diffu- sion rates, which govern creep, are encountered.
Cold-worked austenitic stainless steels have a markedly diminished proportional limit, partic- ularly in compression. This Bauschinger effect, which arises from the easy mobility of disloca- tions, can be eliminated by stress relieving at around 350 °C (660 °F) for 2 h, which provides the thermal energy for dislocation interactions to lock into place. This produces a sharp yield point without premature nonproportional elastic deformation.
Ferritic Stainless Steels
Ferritic stainless steels, from an annealing point of view, must be discussed in two cate- gories. First are the modern, stabilized alloys, which are ferritic at all temperatures. These al- loys behave as interstitial-free (IF) alloys be- cause the interstitial carbon and nitrogen are re- moved from solution as a stable precipitate. In the second category are the older ferritic steels, which have enough austenitizing elements, usu- ally carbon, in solution to cause them to form austenite at what would otherwise be a good an- nealing temperature. This makes them truly quasi-martensitic alloys, and they must be treated accordingly. Table 2, which lists heat- treating temperatures for ferritic stainless alloys, also shows which grades fit into which category. Soaking
Heating of ferritic stainless for hot working is straightforward. Whether stabilized or not, these alloys are heated to the 1000 to 1100 °C (1830 to 2010 °F) range for hot working. The superferritics can be heated to up to 1300 °C (2370 °F). At this temperature, no debilitating phases occur, and ductility is good. The high diffusion rate inherent to the ferritic structure makes homogenization easy. As long as hot working is completed at temperatures above that at which austenite forms, good hot ductility is expected. This is not a concern with IF alloys, which do not form austenite.
Annealing
The IF ferritics do not undergo any phase change with temperature during the course of properly executed heat treatment. The objective of annealing is generally simply to remove the effects of cold work. This is because they do not need to have carbon put into solution and, ex- cept in rare cases, do not have intermetallic
phases that require dissolution. Alloys with high chromium and molybdenum contents can form
σ and/or α', the brittle, ordered body-centered
cubic (bcc) phase, at temperatures below an- nealing temperatures, so rapid cooling is pru- dent when chromium plus molybdenum ex- ceeds 20%.
The driving force for recrystallization in these alloys is limited by the lower stored energy from deformation inherent in the bcc structure. In addition, the pronounced deformation texture leads to annealing responses that are more accu- rately characterized as recovery and grain growth with diminished recrystallization. These alloys retain this texture after annealing, and this characteristic anisotropy is exploited for good drawability. The major concern is to avoid excessive annealed grain size, which greatly re- duces toughness. Anneal at the higher end of the range only if the loss of toughness associated with large grain size is not a concern. Stabiliz- ing anneals are normally unnecessary for stabi- lized ferritics as their high diffusion rates ensure freedom from knife-edge attack due to sensiti- zation from free unbonded carbon combining with chromium at grain boundaries. The stabi- lizing additions of titanium and/or niobium tie up the carbon as stable TiC or NbC, which does not redissolve during annealing.
The interstitial-bearing ferritic stainless steels must be annealed subcritically, or the for- mation of austenite at higher temperatures would make martensite formation on cooling virtually unavoidable. Thus, a typical primary anneal cycle for a typical alloy such as 430 would be nearly 24 h at 750 °C (1380 °F), the majority of which is thermal equilibration of the large coil mass. The actual time at tempera- ture required is less than 1 h. Continuous an- nealing is not practical because the diffusion of carbon is too slow to occur in the dwell time at temperature typical in continuous annealing lines. This cycle also precipitates essentially all the carbon and nitrogen as mixed Cr/Fe car- bides and nitrides and homogenizes chromium content. This necessarily slow process permits subsequent subcritical annealing for mechani- cal properties (to alleviate the effects of cold work) to be done in a few minutes since carbon has been eliminated from solution by the for- mation of fairly stable carbides. Since the ma- terial is generally purchased in the annealed condition, the user need never be concerned with such lengthy anneals.
Table 2 Recommended annealing temperatures for ferritic alloys
Annealing Annealing
Alloy temperature, oC temperature,oF
Stabilized, Cr+Mo<20 409, 439,18 SR 870–925 1600–1700 Unstabilized, Cr+Mo<20 405, 430, 434, 436 705–790 1300–1450 Stabilized, Cr+Mo>20 29-4C, Monit, Seacure, 444 1010–1065 1850–1950 Unstabilized, Cr+Mo>20, 446 760–830 1400–1525
Stress relieving is rarely a concern for any type of ferritic stainless. Unstabilized grades should not be welded, and if they are, full sub- critical annealing is required. Stabilized grades have no need for postweld heat treatment. Low- temperature heat treatment runs the risk of α' formation and is best avoided.
Martensitic Stainless Steels
The martensitic stainless steels resemble the unstabilized ferritic stainless steels described. The martensitic stainless steels form essentially 100% austenite on heating and have very high hardenability, so their ability to be softened by annealing is limited. The traditional martensitic stainless steels are iron/chromium/carbon al- loys, sometimes with a small amount of nickel and/or molybdenum. More recently, alloys have been developed for petroleum applications that contain high copper, nickel, and/or molybde- num and low carbon. The principles of heat treatment of the two alloy categories are the same. The more highly alloyed newer alloys are, in fact, simpler to heat treat because their low carbon and nitrogen levels alleviate the need to temper.
Soaking
Hot working should be carried out in the austenitic range. Temperatures for this are listed in Table 3. Forging and hot working should always be followed by annealing to avoid stress cracking due to the deep hardening of these alloys.
Annealing
Martensitic stainless steels can be annealed by subcritical anneal and sometimes by full an- neal depending on alloy level. If the alloy level is such, as in the nickel-containing grades, that
martensite cannot be avoided by furnace cool- ing from austenitic temperatures, then only sub- critical annealing is feasible. But, even for nickel-free alloys the hardenability is so great that annealing by slow cooling is quite difficult. Martensitic alloys are put into the annealed con- dition for processing before they are quenched and tempered for their final use. Thus, the more economic subcritical anneal is the predominant annealing heat treatment.
The nickel-bearing alloys have such high hardenability that annealing in the critical range cannot produce softening by any practical cool- ing rate, so subcritical annealing is always rec- ommended for these alloys. Nickel reduces the temperature at which austenite is stable as shown in Chapter 9, Fig. 9. Other additions like vanadium, molybdenum, and tungsten promote secondary hardening and tempering resistance, and subcritical annealing of these alloys be-